Copper-deposited carbon fiber aluminum foam sandwich panel performance characteristics compared to conventional aluminum foam sandwich panel

Aluminum foam sandwich (AFS) panels have a combination of structural and functional characteristics such as low density, high specific strength, favorable energy absorption, damping capacity, and electromagnetic shielding. These excellent properties make AFS have great potential in aerospace, transportation, and marine fields. Compared with barely metal foams, the outer metallic plates of AFS can effectively seal the foam and improve its comprehensive mechanical properties, such as tensile and bending strength. Generally, the AFS was fabricated by adhesive bonding method, which entails attaching the aluminum foam to solid metal panels using epoxy resin adhesive.

Although the adhesive bonding method effectively enhances structural damping, it is accompanied by limitations, including challenges in recycling and susceptibility to delamination failure under high-temperature or corrosive conditions. To address these limitations, extensive research efforts have been dedicated to achieving metallurgical bonding between the panel and the core layer.

Currently, the predominant preparation methods for AFS aimed at achieving metallurgical bonding are generally classified into three categories: welding, melt foaming, and packing rolling powder metallurgy method. Among these methods, the packing rolling powder metallurgy method is regarded as superior for controlling pore structure and achieving low-temperature foaming. Our previous work employed this technology to fabricate large-size AFS with better expansion properties and cellular structure for industrial applications.

However, due to the relatively low compressive mechanical properties of the foam core, the improvement of the mechanical strength of the AFS is limited, so further improvement of the strength of the foam core is crucial for the AFS.

To date, the common methods reported to improve the compressive mechanical properties of aluminum foam are as follows: surface treatment, heat treatment, and reinforcement by adding particles, fibers, alloying elements, etc. Among them, short carbon fibers with low density, high elastic modulus, and large aspect ratio have attracted wide attention as promising reinforcing materials. Moreover, surface metallization of carbon fibers, by electroplating or chemical plating such as copper plating or nickel plating, further improves the wettability of the carbon fibers in the aluminum melt and avoids harmful interfacial reactions between the carbon fibers and the aluminum melt. Short carbon fiber is now widely used to reinforce aluminum matrix composites.

Bhav Singhetal. prepared copper-coated carbon fibers reinforced aluminum matrix composites by stir casting method. Muetal. fabricated aluminum matrix composites reinforced with nickel-coated carbon fibers using the two-roll casting method. The results demonstrated a significant enhancement in the tensile strength of the composites with the addition of carbon fibers. However, there has been limited research conducted on aluminum foam reinforced with copper-coated carbon fibers. Caoetal. first prepared copper-coated carbon fibers reinforced aluminum foam through the melt foaming method. The results demonstrate that 0.35 vol% of Cf can stabilize the aluminum foam by preventing liquid film rupture, and the more the amount of fibers added to the foam is more stable. Muetal. showed that the damping properties of aluminum foam prepared by the melt foaming method were significantly improved with the addition of Cf.

Accordingly, it can be concluded that the addition of Cf can be an effective strategy to simultaneously improve the foaming stability and mechanical properties of aluminum foams. Furthermore, there have been no reports on copper-plated carbon fibers reinforced aluminum foam sandwich panels prepared via the powder metallurgy method, which motivates researchers to carry out more in-depth studies in this field.

Numerous studies have been carried out on the nucleation mechanism and stability of aluminum foams with the addition of reinforcing phases. The differences in pore number and pore morphology,

such as size, distribution, and shape, are usually attributed to nucleation mechanisms and stability. Heterogeneous nucleation effects and oxide network stabilization mechanisms in powder metallurgy system has been exhaustively investigated. Quaternary alloy powders consisting of Al–Si–Mg–Cu used in this paper have exhibited exceptional expansion rates and low foaming temperatures, with a solidus temperature of 507 ◦C and a liquidus temperature of 596 ◦C. Nevertheless, there is a lack of research regarding the nucleation mechanism, stability, and compressive mechanical properties of powder compacts containing this particular alloy composition, particularly when incorporating Cf. In previous works, investigations of nucleation and stabilization mechanisms were primarily conducted through non-in situ methods, involving the tomographic analysis of samples that were rapidly cooled following foaming interruption. Unfortunately, submicron bubble nucleationprimarily occurs during precursor heating and expands rapidly within time intervals spanning from 100 ms to 1 s, often presenting challenges for real-time observation. The modern synchrotron radiation X-ray imaging technique possesses high energy and spatiotemporal resolution, allowing for real-time, in situ observation of the nucleation and growth process of aluminum foams at elevated temperatures. Currently, only limited researchers have explored the foaming kinetics of Al–Si–Mg alloy powder by synchrotron radiation. For example, García-Morenoetal. reported significant dynamic phenomena in liquid aluminum foam, including nucleation and growth, bubble rearrangement, liquid retraction, agglomeration, and film rupture. Kamm et al. studied the nucleation and growth process of AlSi8Mg4 aluminum foam by synchrotron radiation. Their findings indicated that hydrogen generated through the decomposition of adsorbates on the metal powder’s surface represents another significant nucleation mechanism. Yet, the interaction among precursor components is intricate, and the dynamics of foaming can be influenced by their composition. For both the Al–Si–Cu–Mg powder system and the Al–Si–Cu–Mg powder system with added Cf, understanding and exploring the initiation and growth of metal foam is crucial for the further structural development of the foam to improve properties for applications.

In this work, Cf/AFS and AFS with metallurgical interfacial bonding were prepared by the packing rolling powder metallurgy method. The bubbles evolution of Cf/AFS and AFS was observed in situ by synchrotron radiation to reveal the mechanism of Cf on bubble nucleation, growth and merger, and stability. In addition, Cf/AFS and AFS were foamed at 620 ◦C for 15 min in a chamber furnace. Subsequently, the pore size distribution, pore microstructure, and compressive mechanical properties of the foamed samples were characterized and tested. Based on the aforementioned studies, the structure and compressive strength of Cf/AFS and AFS were systematically evaluated.

2. Materials and methods

2.1. Raw materials

Six raw materials were used for preparing the AFS and Cf/AFS: Al powder (average size: 45 μm, purity >99.70%), Si powder (average size: 38 μm, purity >99.50%), Cu powder (average size: 38 μm, purity >99.90%), Mg powder (average size: 75 μm, purity >99.70%), TiH2 powder (average size: 45 μm, purity >99.70%), and copper-coated carbon fiber (average size: 1～2 mm, thickness of copper-coated: 1.4 μm).

2.2. Fabrication of Cf/AFS

The schematic of the packing rolling powder metallurgy method is illustrated in Fig. 1. During the mixing stage, the quaternary alloy AlSi6Mg4Cu4 powder was prepared, with powders of Al, Si, Cu, and Mg in proportions of 86 wt%, 6 wt%, 4 wt%, and 4 wt%, respectively. 1 wt% oxidized TiH2 powder (heat treatment in air for 1.5 h at 470 ◦C) was added as a blowing agent which release H2 rapidly at the melt temperature expanding the melt. Additionally, 0.5 wt% of carbon fibers were incorporated as reinforcing phases. Notably, to enhance wettability and prevent harmful reactions like Al4C3 at the fiber-aluminum melt interface during foaming, carbon fibers were copper-coated via electroplating. The electroplating process was described in Refs. Fig. 2 shows the copper-coated carbon fibers obtained by the electroless plating process. The coating has uniform and continuous coverage of the fiber surface (Fig. 2(a)–(b)). Subsequently, the different alloy compositions utilized in this study were prepared by thoroughly mixing aforementioned powders with and without Cf in a three-dimensional mixer for 3 h. The uniformly mixed powder is encapsulated within a sealed cavity, followed by cold and hot rolling procedures.

The cold rolling stage effectively evacuates air from within the cavities and the interstitial spaces between the particles through a series of multiple passes and small depressions. The hot rolling stage ensures that the precursor achieves a relative density exceeding 98% through the co-deformation of the panel and powder, which facilitates the subsequent foaming process. Detailed rolling process parameters are described in Refs. AFS and Cf/AFS foamable precursors are obtained through the aforementioned process. Ultimately, both foamable precursors were heated at 620 ◦C for 15 min to fabricate Cf/AFS and AFS specimens, which were used for pore size analysis, microstructural characterization, and compressive mechanical properties tests.

2.3. Characterization and test methods

2.3.1. Microstructural characterization

The Microstructure morphology of the Cf/AFS and AFS were investigated by field emission scanning electron microscopy (SEM, Zeiss Ultra Plus). Energy dispersive spectroscopy (EDS) was used to identify the composition of various phases in the cell walls of the foamed samples. Longitudinal sections of the samples were obtained by cutting the foamed specimens using the electro-discharging machine (EDM, DK7745). Following this, the specimens’ longitudinal sections were coated with paint and then sanded and polished using 800-grit and 1500-grit sandpaper, respectively. Ultimately, the equivalent cell diameters (Dmean) of the polished specimen cross-section were determined using image analysis software, Image Pro Plus 6.0.

Fig. 1. Schematic of the packing rolling powder metallurgy method.

Fig. 2. SEM images of the (a) copper-coated carbon fibers, and (b) shows a magnified view of the region marked by a red square in (a). Images (c) and (d) are the EDS results of the point A in (b).

Fig. 3. Schematic diagram of the synchrotron radiation experiment setup.

2.3.2. In-situ foaming observations by synchrotron radiation

The schematic diagram of the synchrotron radiation experiment setup is sketched in Fig. 3. The experiments were conducted at beamline 4W1A of the Beijing Synchrotron Radiation Facility (BSRF, Beijing Institute of High Energy Physics, Chinese Academy of Sciences, China). Unlike conventional absorption-contrast imaging, this paper utilizes phase-contrast imaging with the Synchrotron Radiation Light Source. This approach addresses the limitations of traditional absorption imaging, which is ineffective for visualizing weakly absorbing substances. The foaming process of the specimens was performed within a customdesigned foaming furnace, featuring a 15 mm by 15 mm square window positioned perpendicular to the X-ray beam direction. The sample size, as depicted in Fig. 3, measures 15 mm × 15 mm × 2 mm. The sample is placed in a foaming mold consisting of a certain number of ceramic sheets. Throughout the foaming process, X-rays can pass through the sample and are captured by a charge-coupled device (CCD) camera. The chosen X-ray source operated at an intensity of 20 Kev, and a maximum viewing size of 7.4 mm × 7.4 mm was employed to enable comprehensive observation of the entire foaming process.

It is worth noting that the observed samples were positioned in a manner perpendicular to both the rolling direction and the tangential direction. Both the precursors containing 0.5 wt% Cf and those without 0.5 wt% Cf had their side panels removed from both sides of the samples using EDM. The effect of Cf on the foaming characteristics of aluminum foam was investigated. Furthermore, the precursors with different compositions were foamed in the same heating furnace at a preset temperature of 620 ◦C.

2.3.3. Compression mechanical properties test

The compression specimens were cut into cylindrical specimens with a height of 23 mm and a diameter of 20 mm by EDM, and every direction was ensured to include at least 10 complete pores to avoid size effect. Quasi-static compression tests were conducted following the standards ISO13314–2011 and GB/T31930–2015, using an AG-XPLUS electron universal material testing machine witha maximum load capacity of 100 KN. Three samples were tested for each group, and the compressive stress-strain curves in this paper represent the average results obtained from the three tested samples. All tests were performed under displacement control, maintaining a constant cross-head speed of 2 mm/ min (corresponding to an initial strain rate of 10− 3 s− 1) at room temperature.

Fig. 4. Time-evolved synchrotron radiation images of Cf/AFS foaming process: (a) unfoamed precursor at t0 moment of 380 ◦C; (b)–(c) dissolution of elemental copper and type II nucleation during precursor melting; (d)–(g) heterogeneous nucleation and growth processes of bubbles; and (h)–(i) bubbles merging and rupture.

3. Results and discussion

3.1. In situ foaming behavior under synchrotron radiation

3.1.1. Foam evolution of Cf/AFS

Fig. 4 displays a series of synchrotron radiation images depicting the foaming process of the Cf/AFS. The original images of the Cf/AFS (Fig. 4 (a)) were captured when the matrix was in the solid state at 380 ◦C, and this time was set as t0. Compared to the other components in the precursor, Cf appears black within the predominantly grey aluminum-rich matrix. This can be attributed to the high copper content on the fiber surface, as shown in Fig. 2(c)–(d), which has the largest relative atomic mass compared to the other components of the precursor. Moreover, the majority of Cf are observed to be dispersed and separate, rather than clustered together, as exhibited in Fig. 4(a). This implies that with a 0.5 wt% addition of Cf, a 3 h stirring process can effectively attain the desired state of dispersion.

At t0+194s (as seen in Fig. 4(b)), a considerable number of white, nearly spherical bubbles with diameters ranging from 90 μm to 180 μm can be observed in the grey aluminum matrix (marked with the red arrow). The number of bubbles continued to increase within 46 s, as shown in Fig. 4(c). In addition, another noteworthy change taking place in the matrix is that some small bubbles tend to form on the surface of Cf with the initiation of Cu dissolutions (marked with the red dotted box). This phenomenon is likely attributed to type II nucleation occurring within the copper-containing powder compact. It is widely acknowledged that the addition of Cu significantly decreases the melting temperature of alloys. This results in localized melting around individual powder particles in the powder precursor containing copper powder during the heating process, generating a substantial amount of Al–Si–Mg–Cu quaternary eutectic melt at lower temperatures. The quaternary eutectic melt pool tends to be a weak region in the alloy and is more prone to nucleation. As a result, bubbles are more likely to nucleate and grow near the copper-coated on the fiber surface. Furthermore, the low liquid fraction in the matrix at this moment causes gases from the decomposition of TiH2 to easily accumulate in the aluminum matrix. Type II nucleation can effectively avoid the accumulated gas forcing the low liquid fraction matrix apart and propagating cracks through the liquid, resulting in severe loss of the amount of hydrogen.

Fig. 4(d) shows the bubbles initially nucleate and grow on the region of the Cf as the matrix entirely melts, and the bubble locations correspond to the disappearing copper-coated layer (marked with the red dotted box). It is well-established that the addition of secondary phases increases the number of heterogeneous nucleation sites, thereby reducing the energy barrier for the formation of new bubbles within the alloy melt. Meanwhile, nucleation rate is related to the total heterogeneous nucleation area per unit volume. Consequently, it can be observed that the bubbles initially nucleate at the sites where carbon fibers are situated and expand along the length of the carbon fiber. With a further increase in the melt fraction, numerous new bubbles continue to emerge from Fig. 4(e)–4(g). However, the rate of localized merger and rupture of the bubbles is slow, indicating that the bubbles possess high stability. At t0+426s, the number of bubbles increases significantly, and

the precursor volume exhibits large expansion (Fig. 4(h)). Furthermore, the occurrence of some bubble merging (marked with the red arrow), where small bubbles are absorbed by larger ones due to Ostwald maturation, is an inevitable process during the evolution of metallic foams [40]. This maturation phenomenon also results in the appearance of a small number of large and irregular bubbles (marked with the red arrow in Fig. 4(i)).

Fig. 5. Time-evolved images of the foaming process of AFS at different stages: (a) unmelted precursor at t0 moment of 380 ◦C; (b) dissolution of elemental copper and type II nucleation during precursor melting; (c) nucleation and growth processes of bubbles; and (d)–(f) bubbles merging and rupture.

3.1.2. Foam evolution of AFS

Fig. 5 demonstrates a series of synchrotron radiation images illustrating the foaming process of AFS. Similar to Cf/AFS, the initial image for the synchrotron radiation sequence was captured when AFS was in the solid phase at 380 ◦C (see Fig. 5(a)). When the matrix is heated above the solidus temperature, small white bubbles become visible in Fig. 5(b) and (c), resulting from the decomposition of TiH2 and the occurrence of type II nucleation in the Cu-rich region. However, the number of bubbles generated and their distribution density in Fig. 5(c) is notably lower compared to Fig. 4(c). This suggests that Cf/AFS provides more channels for early gas release during nucleation, preventing the excessive for-mation of cracks in aluminum matrix with low liquid fraction.

A fewer number of AFS nucleation sites and a higher probability of abnormal bubble growth compared to Cf/AFS can be observed from Fig. 5(d)–(f). The evolution of AFS bubbles is irregular, and their sizes exhibit wide fluctuations due to the coexistence of both micrometer sized bubbles and abnormally millimeter-sized large bubbles (marked with the red arrow). Moreover, this instability and inhomogeneity during the early nucleation process can result in significant differences in the final pore structure and mechanical properties between AFS and Cf/AFS.

3.2. Macrostructure analysis

To investigate the differences in macroscopic pore morphology and pore diameter between AFS and Cf/AFS, two sandwich structures were prepared under laboratory conditions at a foaming temperature of 620 ◦C for 15 min. The macrostructures and cell size distribution of AFS and Cf/AFS are shown in Fig. 6. The comparison of vertical sections of the two sandwich structures samples reveals that the Cf/AFS foams have finer cells and less defects such as rupture or merging of cells (see Fig. 6 (b)). On the contrary, as shown in Fig. 6(a), the AFS exhibits poor pore homogeneity with the presence of some abnormally large pores. Meanwhile, the equivalent pore diameter for AFS and Cf/AFS were 2.9 ± 0.2 mm and 1.9 ± 0.1 mm, respectively, suggesting that the pore diameter was further refined by the addition of Cf (see Fig. 6(c) and (d)).

The evolution of the pore structure is influenced by two competitive mechanisms: the decomposition of TiH2 and the merging and rupture of the foam . The decomposition of TiH2 contributes to the increase in the number of nucleation, porosity, and expansion rate, while the merging and rupture of bubbles decreases the number of pores and increases the diameter of the pores. The reduction of bubble merging and rupture events contributes to obtaining a better cellular structure. Hence, the high stability of the foam is necessary to maintain the homogeneity of the pore structure during the foaming process. For Cf/AFS, the addition of Cf greatly improves the wettability of the fibers with the aluminum melt. Similar to the reported Cf reinforced aluminum foam [44] prepared by the melt foaming method, the addition of Cf stabilizes the aluminum foam by preventing cell wall rupture and reducing agglomeration. Meanwhile, Cf can create a net-like structure in the cell wall, which generates a separation force and prevents the rupture of bubbles, thus, improving the cell structure, foam stability, and uniformity of the pore distribution.

3.3. Microstructural characterization

Fig. 7(a) shows the microstructure of the cell wall after foaming with AFS. As reported for AlSi6Cu4Mg4 alloy foams, eutectic Al–Si, Mg2Si, Al2Cu, and Al5Cu2Mg8Si6, are present in the alloy, which is in agreement with the findings obtained from EDS analysis in Fig. 7(b). These brittle phases decrease the mechanical properties of the aluminum foam by weakening the pore wall or acting as a source of cracks during deformation. Therefore, enhancing the mechanical properties of the pore wall through the addition of reinforcing phases is a viable and effective strategy.

For Cf/AFS, it can be observed from Fig. 7(c) and (d) that the Cf is distributed inside the cell wall or at the edge of the cell wall, indicating Cf has excellent wettability with the aluminum melt. Furthermore, it is noticeable (refer to Fig. 7(d)) that a minor amount of the precipitated phase Al2Cu forms at the interface between the fibers and the aluminum melt. This phenomenon contributes to achieving a strong bond between the fibers and the matrix. The progression of this diffusion reaction can be intuitively observed in Fig. 4(a)–(c). Note that, the precipitated phase does not completely surround the carbon fibers, and the state and amount of this distribution are beneficial for Cf/AFS. Based on the study of Lv et al, on Cf reinforced aluminum matrix composite materials, it was demonstrated that the appropriate amount of precipitated phase can influence and regulate the interfacial bond strength.

Fig. 6. Vertical sections of representative sample (a) AFS and (b) Cf/AFS. The equivalent diameter versus area fraction plots corresponding to (a) and (b) are depicted in (c) and (d), respectively. The solid continuous line in (c) and (d) represents the log-normal fit. Regression (R2 ) represents the goodness of fit.

Fig. 7. SEM images (a) and (b) show the microstructures corresponding to Fig. 6(a) AFS and Fig. 6(b) Cf/AFS, respectively, furthermore, Fig. 7(b) and (d) show localized magnifications of Fig. 7(a) and (c), respectively.

3.4. Mechanism of Cf stabilizing Al–Si–Mg–Cu foam

Based on the above-mentioned analysis, a schematic representation of the foaming mechanism for Cf/AFS and AFS is summarized in Fig. 8. The significant differences between Cf/AFS and AFS in the nucleation stage, bubble growth, and merging stage, and the gradual solidification stage, as illustrated in Fig. 8(a) and (b).

Fig. 8. Schematic diagram of Cf/AFS and AFS foaming behavior: (a) Cf/AFS; and (b) AFS.

In the nucleation stage of AlSi6Cu4Mg4 alloy foams, bubbles tend to nucleate within the liquid quaternary eutectic surrounding the copper particles, rather than nucleate at the location of the TiH2 particles. This type of nucleation in the weakest region of the precursor (type II nucleation) provides more escape channels for the gases released by TiH2 decomposition, avoiding the accumulation of gases and causing cracks to propagate in the aluminum melt. The addition of copper-coated carbon fibers apparently further creates more channels for releasing H2. However, since the aluminum melt has no sufficient liquid fraction at this stage, it is difficult for such small bubbles to have a sufficient source of H2 to continue to grow. As a result, these small bubbles rupture within the melt as the temperature rises. When the precursor is completely melted, the H2 released from the violent dehydrogenation reaction of TiH2 will nucleate and grow rapidly based on the nucleation point. The Cf provides numerous heterogeneous nucleation sites for the generation of bubbles, which greatly increase the nucleation rates. Simultaneously, the presence of Cf effectively inhibits the contact and merger of the primary bubbles and avoids the formation of anomalously large bubbles.

During the growth and merger stage of the bubbles, the addition of Cf significantly improves the foam stability. The reason is that Cf and aluminum melt have good wettability, and the fibers can be easily trapped between the two gas-liquid interfaces to form a network structure (as shown in Fig. 7(c)). As the liquid film is thinned, the fibers can generate a separation pressure to resist suction from the plateau boundary, thus preventing the liquid film from rupturing. So Cf/AFS needs longer foaming time compared to AFS, with higher stability that can maintain the uniformity of the pore structure during the evolution of the foam.

During the slow solidification stage, bubbles in Cf/AFS no longer grow and the fibers are embedded within the cell wall (marked with the red dotted box). This retention of fibers prevents the absorption of small bubbles by larger ones, as seen in Ostwald maturation, and inhibits cell wall rupture due to solidification expansion. As a result, Cf/AFS exhibits a more homogeneous pore structure and fewer pore defects compared to AFS.

3.5. Mechanical properties

The compressive stress-strain curves for Cf/AFS and AFS with almost the same porosity (83.2% and 81.4%, respectively) are presented in Fig. 9(a). The compressive stress-strain curve displays a distinct threestage characteristic, comprising the linear elastic region, plateau region, and densification region. The compressive yield strength of AFS and Cf/AFS were 8.44 MPa and 11.87 MPa, respectively. The compressive yield strength of Cf/AFS was increased by 40.6% compared to AFS. Meanwhile, the energy absorption capacity of Cf/AFS is better than that of AFS at the same strain, as shown in Fig. 9(b). The energy absorption capacities of AFS and Cf/AFS are about 3.3 MJ/m3 and 6.1 MJ/m3, respectively. The energy absorption capacity of Cf/AFS is improved by 84.8% compared with AFS.

Fig. 9. Compressive stress-train curves and energy absorption capacity curves of Cf/AFS and AFS: (a) compressive stress-train curves; (b) energy absorption capacity curves.

The incorporation of Cf significantly enhanced the mechanical properties of the foam core. It is well known that the mechanical properties of aluminum foam are closely related to the pore structure and strut mechanical properties. Macrostructural analysis indicated that the addition of Cf refined the pore size, improved pore distribution, and reduced pore defects. Similar to the findings of Sasikumar et al. and Yangetal. small pore diameter and narrow pore distribution are more beneficial to enhance the compressive strength and energy absorption efficiency. Moreover, the effect of microstructure on the mechanical properties of the foam is mainly due to the Cf embedded in the pore walls of the foam. The diffusion and dissolution of the copper-coated of the carbon fibers in the aluminum matrix can achieve a firmer interfacial bond between the fibers and the aluminum matrix, which contributes to load transfer. The high modulus carbon fiber stapled in the pore wall can limit the deformation of the pore wall under compression load and improve the load-bearing capacity. All these factors favor Cf/AFS to achieve higher yield strength and energy absorption capacity compared to AFS.

4. Conclusions

The Cf/AFS and AFS were prepared by the packing rolling powder metallurgy method in this work. The foaming behavior of Cf/AFS and AFS was dynamically observed through synchrotron radiation to

investigate the influence of Cf on the bubble nucleation and growth. In addition, Cf on the pore morphology, pore distribution, and compression properties of aluminum foams was thoroughly analyzed. The main conclusions are as follows:

(1) Synchrotron radiation in situ observations reveal that both Cf/ AFS and AFS foaming processes exhibit three distinct stages: bubble nucleation, bubble growth and merger, and solidification.

In comparison to AFS, the foaming process of Cf/AFS demonstrates higher stability and bubble uniformity. AFS is more prone to forming abnormally large and coarse pores during the pore evolution process. Additionally, the disparity is attributed to the heterogeneous nucleation effect of Cf, which prevents bubble wall rupture and reduces coalescence while increasing the nucleation rate.

(2) The addition of 0.5 wt.% Cf refines the equivalent pore diameter from 2.9 ± 0.2 mm to 1.9 ± 0.1 mm, significantly reducing coarse pore defects. Simultaneously, the microstructure reveals good wettability of Cf, which is distributed along the pore wall and gas-solid interface, thereby enhancing foam stability.

(3) Compared with AFS, the compressive yield strength and energy absorption capacity of the Cf/AFS increased by approximately 40.6% and 84.8%, respectively. The findings of this study can provide ideas for the reinforcement of aluminum foam sandwich panels with metallurgical bonding interfaces that have the potential for engineering applications in the field of automotive and aerospace industries.

Article from Journal of Materials Research and Technology